Method of reducing the formation of primary platelet-shaped beta-phase in iron containing alSi-alloys, in particular in Al-Si-Mn-Fe alloys

ABSTRACT

The present invention is a method for producing an iron-containing hypoeutectic alloy free from primary platelet-shaped beta-phase of the Al 5 FeSi in the solidified structure by the steps (a) providing an iron-containing aluminum alloy having a composition within the following limits, in weight percent, 6-10% Si, 0.05-1.0% Mn, 0.4-2% Fe, at least one of 1) 0.01-0.8% Ti and/or Zr 2) 0.005-0.5% Sr and/or Na and/or Ba, 0-6.0% Cu, 0-2.0% Cr, 0-2.0% Mg, 0-6.0% Zn, 0-0.1 % B balance aluminum (b) controlling and regulating precipitation path during solidification such that the precipitation of Fe containing intermetallic phases starts with the precipitation of the hexagonal phase of the Al 8 Fe 2 Si by (b1) controlling the condition of crystallization by addition of one or more of Fe, Ti, Zr, Sr, Na and Ba within the limits specified in step (a) and (b2) identifying the phases or morphology of the phases that precipitates during the solidification and correct the addition one or more times in order to obtain desired precipitation path and (c) solidifying the alloy at the desired solidification rate.

The present invention relates to a method of producing iron-containingAl-alloys having improved mechanical properties, in particular improvedfatigue strength, by controlling the morpholgy of the iron containingintermetallic precipitates.

BACKGROUND OF THE INVENTION

Iron is known to be the most common and at the same time mostdetrimental impurity in aluminium alloys since it causes hard andbrittle iron-rich intermetallic phases to precipitate duringsoidification. The most detrimental phase in the microstructure is thebeta-phase of the Al₅FeSi-type because it is platlet-shape. Since thedetrimental effect increases with increasing volume fraction of thebeta-phase much interest has focused on the possibilites of reducing theformation of said phase, as recently reviewed by P. N. Crepeau in the1995 AFS Casting Congress, Kansas City, Mo., 23-26 April 1995.

The problem related to iron contamination of alumninium alloys is ofgreat economical interest since 85% of all foundry allous are producedfrom scrap, the recycling rate is ever increasing (already higher than72%) and the service life of aluminum is relatively short (of about 14years). As a result thereof, the iron content in aluminium scrapcontinouosly increases since iron cannot be economically removed fromaluminium. Dilution is the only practical method to reduce the ironcontent and the cost of aluminium is known to be inversely related toits Fe content. On the other hand, iron is deliberately added in anamount of 0.6-2% to a number of die-casting alloys, eg BS 1490: LM5,LM9, LM20 and LM24. Moreover, due to the low diffusivity of iron insolid aluminium there exist no practical possibility to reduce thedeleterious effect of the iron containing precipitates by a heattreatment.

Iron has a large solubilty in liquid aluminium but a very low solubiltyin solid aluminium. Since the partition ratio for Fe is quite low, ironwill segregate during solidification and cause beta-phase to form alsoat relatively low iron contents as shown by Bäckerud et al in“Solidification Characteristics of Aluminium Alloys”, Vol. 2,AFS/Skanaluminium, 1990. In said book the composition and morphology ofiron containing intermetalic phases are detailed in relation to theAl-Fe-Mn-Si system.

The two main types occuring in Al-Si foundry alloys are the Al₅FeSi-typephase and the Al₁₅Fe₃Si₂-type phase. Moreover, a phase of theAl₈Fe₂Si-type may form. These intermetallic phases need not bestoichiometric phases, they may have some variation in composition andalso include additional elements such as Mn and Cu. In particularAl₁₅Fe₃Si₂ may contain substantial amounts of Mn and Cu and couldtherefore be represented by the formula (Al,Cu)₁₅(Fe,Mn)₃Si₂.

However, for typing reasons the simplified formulas Al₁₅Fe₃Si₂, Al₈Fe₂Siand Al₅FeSi are preferred in the following. Accordingly, it is to beunderstood that compositional and stoichoimetrical deviations of thephases at issue are covered by the simplified formulas.

The Al₅FeSi-type phase, or beta-phase, has a monoclinic crystalstructure, a plate like morphology and is brittle. The platlets may havean extension of several millimeters and appear as needles inmicrographic sections.

The Al₈Fe₂Si-type phase has a hexagonal crystal structure and dependingon the precipitation conditions this phase may have a faceted,spheroidal or dendritic morphology.

The Al₁₅Fe₃Si₂-type phase (often named alpha-phase), has a cubic crystalstructure and a compact morphology, mainly of the chinese script form.

In the Al-Fe-Mn-Si system these three phases have been represented inthe Si-FeAl₃-MnAl₆-equilibrium phase diagram as described by Mondolfo,FIG. 1. It may be noted that the Al₁₅Fe₃Si₂-type intermetallic isdenoted (Fe,Mn)₃Si₂Al₁₅ in this figure. Point A represents thecomposition of a foundry alloy of the conventional A380-type and it canbe seen that its original composition lies within the (Fe,Mn)₃Si₂Al₁₅area. The solidification of such an alloy typically starts with theprecipitation of aluminium dendrites and, in course of thesolidifcation, the interdendritic liquid becomes sucessivley enriched iniron and silicon. As a result, the Al₁₅Fe₃Si₂-type intermetallic phasestarts to precipitate (represented as(Fe,Mn)₃Si₂Al₁₅ in this diagram).Fe and Mn are consumed due to this reaction. The liquid moves towardsthe Al₅FeSi-area and starts to co-precipitate large platelets ofAl₅FeSi-type phase until the liquid composition reaches the eutecticcomposition at point M in the phase diagram where the main eutecticreaction take place. For further details on the solidification ofcommersial aluminium foundry alloys, reference is given to Bäckerud etal, “Solidification Characteristics of Aluminium Alloys”, Vol. 2,Foundry alloys, AFS/Skanaluminium, 1990.

As already pointed out, the primary platelet-shaped beta-phase of theAl₅FeSi-type is the most detrimental iron containing intermetallic phasein aluminium alloys because of its morphology. The large beta-phaseplatelets have been reported to decrease: ductility, elongation, impactstrength, tensile strenght, dynamic fracture thoughness and impactthoughness. The effect has been attributed to: easier void formation,cracking of the platelets and microporosity caused by the largebeta-phase platelets. In addition, the coarse beta-phase platelets havebeen reported to infer with feeding and castability and thereby increasethe porosity. The perhaps most important effect of the platelets formany industrial applications is that they give rise to microporositywhich is the most likely source of crack initiation.

In summary, it can be concluded that increased Fe may result inunexpected formation of the deleterios platelet-shaped beta-phase. Thebeta-phase forms above a critical iron content, causing the mechanicalproperties to decrease drastically.

Accordingly, in the prior art much work has been directed to thepossibilites of avoiding the formation of beta-phase.

Prior art methods for reducing the formation of beta-phase can begrouped into the following four classes:

1. Control of Fe-content.

2. Physical removal of Fe.

3. Chemical neutralization.

4. Thermal interaction

The first method is based on careful control and selection of the rawmaterials used (ie low-Fe scrap) or dilution with pure primaryaluminium. This method is very costly and restricts the use of recycledaluminium.

The second method relates to sweat melting and sedimentation of ironrich intermetallic phases by the so called sludge. However, both methodsresult in considerable aluminium losses (about 10%) and are thereforeeconomically unacceptable.

Chemical neutralization is, so far, the most used technique. Chemicalneutralization aims at inhibit the platelet morphology by promoting theprecipitation of the Al₁₅Fe₃Si₂-type phase which has a chinese scriptmorphology by the addition of a neutralizing element. In the past, mostwork has been directed to use of the elements Mn, Cr, Co and Be.However, these additions have only been sucessful to a limited extent.Mn is the most frequently used element and it is common to specify %Mn>0.5(% Fe). However, the amount of Mn needed to neutralize Fe is notwell established and beta-phase platelets may occur even when % Mn>% Fe.This method can be used to suppress the formation of beta-phase.However, it is to be noted that the total amount of iron containingintermetallic particles increases with increasing amount of manganeseadded. Creapeau has estimated that 3.3 vol. % intermetallic form foreach weight percent of total (% Fe+% Mn+% Cr) with a correspondingdecrease in ductility. In addition, large amounts of Mn are costly.Chromium and Co have been been reported to act similar as Mn and bothelements suffer from the same drawbacks as Mn. Beryllium works inanother way in that it combines with iron to form Al₄Fe₂Be₅, butadditions >0.4% Be are required which causes high costs in addition tothe safety problems related to the handling of Be since it is a toxicelement.

The last method—thermal interaction—can be performed in two ways.Firstly, by overheating the melt prior to casting in order to reducenucleating particles that form the detrimental phases. However, hydrogenand oxide contents increases, process time is consumed and costs areincurred. The second possibility is to increase the cooling rate in thecombination with an addition of Mn. By increasing the cooling rate theamount of Mn needed decreases somewhat. Although this technique limitsthe drawbacks of the chemical neutralization by Mn it may be hard orimpossible to put into practice in commercial foundry production, inparticular for conventional casting in sand moulds and permanent mouldswith sand cores.

Accordingly, the object of this invention is to propose an alternativemethod to avoid the formation of the deleterious plate like beta-phasein iron containing aluminium alloys. In particular, it is an object topropose a method which does not suffer from the above mentionedproblems.

SUMMARY OF THE INVENTION

In accordance with the invention, this object is accomplished by thefeatures of claim 1. Preferred embodiments of the method are shown independent claims 2 to 10. Claim 11 defines the use of thermal analysisfor controlling the morphology of iron containing intermetallicprecipitates in iron containing aluminium alloys according to claim 1and claim 12 defines a preferred embodiment of claim 11.

The method according to this invention is based on the finding that theprecipitation of platelet-shaped beta-phase of the Al₅FeSi-type can besuppressed by a primary precipitation of the hexagonal Al₈Fe₂Si-typephase. The presence of said Al₈Fe₂Si-type phase result in that whenbeta-phase precipitates it will not develop the commonplatlet-morphology but rather nucleate on and cover the Al₈Fe₂Si-typephase which in turn has a less harmful morphology.

The method of the invention has a number of advantages. Since theprecipitation path during solidification can be controlled to avoid theformation of beta-phase platlets, the iron content need not bedecreased. In apparent contrast to conventional practice, allowable ironcontents may even be increased since iron can influence positively onthe precipitation of Al₈Fe₂Si-type phase. As a result, cheaper rawmaterial can be used. Due to the fact that Mn-additions can be avoided,alloy costs are saved and ductility increases as far as the total amountof iron containing intermetallic particles is reduced.

BRIEF DESCRIPTION OF THE DRAWINGS

The invention will now be described in relation to some examples andwith reference to the accompanying figures in which:

FIG. 1 is a part of the Al-Fe-Mn-Si system as described by Mondolfo. Itdiscloses the Si-FeAl₃-MnAl₆-equilibrium phase diagram.

FIG. 2 shows principally the result of a thermal analysis of analuminium A380-type alloy, wherein the solidification rate (relativerate of phase transformation)(dfs/dt) has been represented as a functionof the fraction solid (fs).

FIG. 3 shows principally the result of a thermal analysis of a boronalloyed A380-type alloy represented in same way as in FIG. 2.

FIG. 3a discloses the result prior to regulation of the crystallizationpath and FIG. 3b shows the result after addition of the precipitationregulating agents(0.15% Ti and 0.02% Sr).

DETAILED DESCRIPTION OF PREFERRED EMBODIMENTS

Thermal analysis was performed for an A380 aluminium alloy with andwithout the addition of a crystallization modifying agent. The analysisof the base alloy is given in Table 1.

TABLE 1 Chemical composition of the base alloy A380 (in weight %). Si9.04 Mn 0.29 Fe 0.95 Cu 3.1 Cr 0.06 Mg 0.04 Zn 2.3 Ti 0.04 Ni 0.12 Sr<0.01

balance Al, apart from impurities.

Sample A represents the base alloy and sample B an alloy to which Ti andSr were added in amounts of 0.1% and 0.04%, respectively. Ti was addedto the melt in the form of an Al-5% Ti-0.6% B alloy and Sr in the formof an Al-10% Sr alloy, the former gave rise to a B content of 0.012% inthe melt. The position of both alloys lies within the (Fe,Mn)₃Si₂Al₁₅area in the Si-FeAl₃-MnAl₆-equilibrium phase diagram and can berepresented by point A in FIG. 1.

About 1 kg of the alloy was melted in a resistance furnace and kept at800 C. Additions were made and the melt was held for 25 minutes at thistemperature. Thereafter the solidification process was investigated bythermal analysis as described by Bäckerud et al in “SolidificationCharacteristics of Aluminium Alloys”, AFS/Skanaluminium, Vol. 1, 1986.The graphite crucible was preheated to 800 C., filled with the melt,placed on a fibrefrax felt, covered with a fibrefrax lid and allowed tocool freely, which led to a cooling rate of approximately 1K/s. Sampleswere taken 10 mm above the bottom of the crucible for metallographicexamination.

In order to examine the nucleation and growth process of the ironcontaining intermetallic phases, specimens were also quenched in waterat specific solidification times.

The solidification process was analysed by conventional thermal analysisas described in the reference given above. Thermal analysis data wascollected in a computer in order to calculate rate of solidification(dfs/dt) and fraction solid (fs) versus time (t). The solidificationprocess was represented by plotting the solidification rate (relativerate of phase transformation)(dfs/dt)as a function of the fraction solid(fs). Curve A (FIG. 2) is from the solidification of the base alloy andcurve B is that of sample B,(0.1% Ti and 0.04% Sr added).

The solidification of the base alloy, curve A, follows the scheme:

Reaction 1 Development of dendritic network

Reaktion 2 Precipitation of AlMnFe containing phases

Reaction 3 Main eutectic reaction

Reaction 4 Formation of complex eutectic phases

The metallographic examiniation of the microstructure of sample Arevealed both beta-phase of the Al₅FeSi-type and Al₁₅Fe₃Si₂-type phaseas iron containing intermetallic phases. In the polished section theplatelet-like beta-phase appeared as large needles and theAl₁₅Fe₃Si₂-type phase as chinese script. The solidification of sample Acan be described in the following manner in relation to FIG. 1, wherepoint A represents the composition of the alloy: First aluminiumdendrites are precipitated and thereafter Al₁₅Fe₃Si₂ starts topricipitate. Mn and Fe are then consumed and point A moves towards theAl₅FeSi area. As a result Al₅FeSi (beta phase) starts to precipitateshortly after the Al₁₅Fe₃Si₂-phase. In FIG. 2 the preciptation ofprimary aluminium is represented by R1 and the precipitation of theintermetallic phases are represented by the two peaks in the R2 area.

The solidification of sample B followed curve B in FIG. 2. In this caseit is to be noted that no peak for reaction 2 could be observed and thatreaction 3 was postponed. A detailed analysis of the data collectedduring the thermal analysis showed that by the additions made to sampleB the liquidus temperature rose about 6 K (the liquidus line KM in FIG.1 moves towards the Al₁₅Fe₃Si₂-area) and the main eutectic reaction waspostponed and occured at a lower temperature. This favours point A to bein or closer to the Al₈Fe₂Si-area. As a result, the fraction solid (fs)at start of the main eutectic reaction (reaction 3) was increased and ina polished section of this sample neither beta-phase of the Al₅FeSi-typenor Al₁₅Fe₃Si₂-phase could be identified. The iron intermetallic phaseprecipitated was identified to be the hexagonal Al₈Fe₂Si-type phasewhich occured as small, mainly faceted, particles. Quenching experimentsshowed that Al₈Fe₂Si-type particles started to precipitate at nearly thesame time as the precipitation of dendritic aluminium. This facetedphase was found to decrease in size and change its morphology fromfaceted to spheroidal with increasing cooling rate. At higher coolingrates, the faceted particles became rather small and homogeneouslydistributed.

All thermodynamic and kinetic factors influencing the formation of ironcontaining intermetallic phases are not known in detail. However, it isthought that the addition of one ore more regulating agents, made inaccordance with this invention to regulate the condition ofcrystallization, acts in one or more of the following ways on theformation of the Al₈Fe₂Si-type phase:

1. Increase in liquidus temperature (eg Ti, Zr).

2. Decrease of the eutectic temperature (eg Sr).

3. Displacement of the starting point in the phase diagram (Fe).

4. Inocculation of the Al₈Fe₂Si-type phase.

The first two points have already been discussed in relation to thesolidification of sample B.

The third mechanism is mainly related to the iron content of thestarting alloy. The iron content infuences the solidfication path in twoways; firstly, the starting point in the Si-FeAl₃-MnAl₆-equilibriumphase diagram is moved towards the iron rich corner of the phase diagramand, secondly, the residual interdendritic melt will enrich more heavilyin iron due to segregation. As a result thereof the melt will firstreach the Al₈Fe₂Si area and cause Al₈Fe₂Si-type phase to precipitate.Finally, it is plausible that complex boride phases form in the melt, egas a result of the use of master alloys for alloying and/or grainrefining purposes. These master alloys often contain borides which, inturn, are known to react with other elements in the melt (such as Sr,Ca, Ni and Cu) to form mixed boride phases. As an example, if Sr ispresent in the melt it will react with the boride particles AlB₂ or TiB₂to form mixed borides having increased cell parameters as compared tothe pure AIB₂ or TiB₂. As a result thereof, the misfit between thehexagonal Al₈Fe₂Si-type phase and the hexagonal borides will decreaseand, hence, favour the nucleation of Al₈Fe₂Si-type phase on the mixedborides.

However, the most important finding is that the precipitation of theplatlet-shaped beta-phase of the Al₅FeSi-type can be suppressed by aprimary precipitation of the hexagonal Al₈Fe₂Si-type phase. It isthought that the precipitation of beta-phase is not inhibited by thepresence of said Al₈Fe₂Si-type phase but that the beta phase cannotdevelop the common platlet morphology since it will nucleate andprecipitate on the Al₈Fe₂Si-type phase. Accordingly, the iron containingintermetallics formed must be supposed to have a core of the hexagonalAl₈Fe₂Si-type phase covered with a layer of the monoclinic beta-phase ofthe Al₅FeSi-type. Since the morphology of these “duplex” intermetallicparticles is governed by the Al₈Fe₂Si-type phase no platlets are formedand the porosity in the solidified structure will be a considerablydecreased. Consequently, the mechanical properties of the final productwill improve, in particular the fatigue strength.

The use of thermal analysis for controlling the morphology is furtherexemplified in relation to sample C which is a boron alloyed (0.1% B)A380-type alloy. A sample of this alloy was taken and analysed bythermal analysis in the same manner as previously described. Byanalysing the curve of the thermal analysis, FIG. 3a, the precipitationof beta-phase could easily be determined and it could also be determinedthat the precipitation started early (ie at a low fs). In order toregulate the precipitation path during solidification such that theprecipitation of the iron containing intermetallic phases starts withthe precipitation of the hexagonal phase of the Al₈Fe₂Si-type aregulating agent was added to the melt in an amount of 0.15% Ti and0.02% Sr. The precipitation path during solidification wasreinvestigated by thermal analysis, FIG. 3b, the absence of the R2-peakand, hence, primary beta-phase is apparent. The melt was then subjectedto casting.

Metallographic samples were taken from both samples as well as from thefinal product and examined by standard metallographic techniques. In thepolished section of the uncorrected sample C, large and long needles ofbeta-phase was observed. However, the structure of the sample examinedafter correction as well as that of the final product no needles ofbeta-phase were observed. The iron containing intermetallic phaseprecipitated appeared as a large number of small faceted particles astypical for the Al₈Fe₂Si-type phase.

Although, thermal analysis is a preferred method to investigate thesolidification path and to identify the precipitation of beta-phaseother methods may be used depending on local factors such as: productionprogram, time limitations and prevailing facilities. From the examplesgiven above it is apparent that the phases precipitated and theirmorphology can be identified by conventional metallo-graphic examinationof a solidified sample. Accordingly, by analysing the structure of asample solidified at a desired solidification rate, it would be possibleto examine the mor-phology of the precipitated phases and thereby toidentify the precence of beta-phase in the structure. The conditions ofcrystallization could then be corrected by addition of one or more ofthe modifying agents Fe, Ti, Zr, Sr, Na and Ba one or more times, ifnecessary, in order to obtain the desired precipitation path. However,this controlling method is deemed to take longer time than thermalanalysis. Alternatively, the chemical analysis might be used tocalculate the activities of the elements in the melt, the position ofthe melt in the actual phase diagram, the segregation duringsolidification and so forth. These data could then be used, alone or incombination with an expert system, for calculation of the solidificationpath of the alloy. In addition, additions necessary to ensure that theprecipitation of the iron containing intermetallic phases starts withthe precipitation of the hexagonal phase of the Al₈Fe₂Si-type couldpossibly be calculated for the desired solidification rate. However, atpresent no such system is fully developed to suit foundry practice.

What is claimed is:
 1. A method for producing an iron containinghypoeutectic aluminium alloy free from primary platelet-shapedbeta-phase of the Al₅FeSi-type in the solidified structure by the stepsof a) providing an iron containing aluminium alloy having a compositionwithin the following limits in weight %: Si 6-10 Mn 0.05-1.0 Fe 0.4-2.0at least one of 1) Ti and/or Zr 0.01-0.8 2) Sr and/or Na and/or Ba)0.005-0.5 optional one or more of Cu 0-6.0 Cr 0-2.0 Mg 0-2.0 Zn 0-6.0 B0-0.1 balance Al apart from impurities, b) controlling and regulatingthe precipitation path during solidification such that the precipitationof Fe containing intermetallic phases starts with the precipitation ofthe hexagonal phase of the Al₈Fe₂Si-type by b1) regulating the conditionof crystallization by addition of one or more of Fe, Ti, Zr, Sr, Na andBa within the limits specified in step a) and b2) identifying the phasesand/or the morphology of the phases that precipitate during thesolidification and, if necessary, correct the addition one or more timesin order to obtain the desired precipitation path, and c) solidifyingthe alloy at the desired solidification rate.
 2. A method according toclaim 1 wherein the identification of the phases and/or the morphologyof the phases that pre-cipitates during the solidification is performedby at least one of thermal analysis, metallographic method and numericalcalculation.
 3. A method according to claim 1 wherein the condition ofcrystallization in step b1) is per-formed by the addition of Ti.
 4. Amethod according to claim 1 wherein the condition of crystallization instep b1) is per-formed by the combined addition of Ti and Sr.
 5. Amethod according to claim 1 wherein the condition of crystallization instep b1) is per-formed by the addition of Fe.
 6. A method according toclaim 1 wherein the solidifcation rate is <150 K/s.
 7. A methodaccording to claim 1 wherein the composition of the liquid alloy lieswithin the (Fe,Mn)₃Si₂Al₁₅-area in the Si-FeAl₃-MnAl₆-equilibrium phasediagram.
 8. A method according to claim 1 wherein the aluminium alloyhas a composition within the following limits in weight %: Si 7-10 Mn0.15-0.5 Fe 0.6-1.5 Cu3-5.
 9. A method according to claim 1 wherein thealuminium alloy has a composition within the following limits in weight%: Si 8.5-9.5 Mn 0.2-0.4 Fe 0.8-1.2 Cu3.0-3.4
 10. A method according toclaim 1 wherein the element or elements regulating the condition ofcrystallization is added in the form of a master alloy.
 11. A methodaccording to claim 1 characterized in that the phases and/or themorphology of the phases that precipitate during the solidification isidentified by using thermal analysis.
 12. A method according to claim 11wherein the data of the thermal analysis is used for controlling andregulating the preci-pitation path during solidification such that theprecipi-tation of Fe containing intermetallic phases starts with theprecipitation of the hexagonal phase of the Al₈Fe₂Si-type.
 13. A methodaccording to claim 3 wherein the amount of Ti added is 0.1-0.3% Ti. 14.A method according to claim 3 wherein the amount of titanium addition is0.15 to 0.25% Ti.
 15. A method according to claim 4 wherein the amountof titanium added is 0.1-0.3% Ti and the amount of strontium added is0.005-0.03% Sr.
 16. A method according to claim 4 wherein the amount oftitanium added is 0.15-0.25% Ti and the amount of strontium added is0.01-0.02% Sr.
 17. A method according to claim 5 wherein the amount ofiron added is 0.5-0.15% Fe.
 18. A method according to claim 5 whereinthe amount of iron added is 0.5-1.0% Fe.
 19. A method according to claim6 wherein the solidification rate is <100 Ks.
 20. A method according toclaim 6 wherein the solidification rate is <20 Ks.
 21. A methodaccording to claim 10 wherein said master alloy contains particles witha hexagonal structure.
 22. A method according to claim 10 wherein saidmaster alloy contains a nucleating agent for the Al₈FeSi₂ phase.
 23. Aniron-containing hypoeutectic aluminum-silicon alloy free fromplatelet-shaped beta-phase of the Al₅FeSi-type having a compositionwithin the following limits in weight percent: Si 6-10 Mn 0.05-1.0 Fe0.4-2.0 at least one of 1) Ti and/or Zr 0.01-0.8 2) Sr and/or (Na and/orBa) 0.005-0.5 optionally one or more of Cu 0-6.0 Cr 0-2.0 Mg 0-6.0 Zn0-6.0 B 0-0.1 balance Al apart from impurities, and containing ahexagonal phase of the Al₁ ₈FeSi₂ type as the primary precipitatedFe-containing intermetallic phase.
 24. An alloy according to claim 23having a composition within the following limits in weight percent: Si7-10 Mn 0.15-0.5 Fe 0.6-1.5 Cu 3-5.
 25. An alloy according to claim 23having a composition within the following limits in weight percent: Si8.5-9.5 Mn 0.2-0.4 Fe 0.8-1.2 Cu 3.0-3.4.